Ni-Based Alloy Product and Method for Producing Same, and Ni-Based Alloy Member and Method for Producing Same

ABSTRACT

There are provided: an Ni-based alloy member including a γ′ phase precipitation with 36 to 60 volume % and exhibiting a high durable temperature and good cold workability; a method for producing the member; an Ni-based alloy product to be used as a precursor of the member; and a method for producing the product. The Ni-based alloy product has a two-phase structure composed of a γ phase and a γ′ phase being incoherent to the γ phase, the incoherent γ′ phase being present at a ratio of 20 volume % or higher. The Ni-based alloy member produced by cold working the Ni-based alloy product and subsequently by conducting heat treatment comprises a γ phase and a γ′ phase being coherent to the γ phase, the coherent γ′ phase being present at a ratio of 36 to 60 volume %, and has a predetermined shape.

CROSS REFERENCE TO RELATED APPLICATIONS

This application is a Divisional of U.S. application Ser. No.14/905,075, filed Jan. 14, 2016, which claims priority fromInternational Patent Application No. PCT/JP2013/069367 filed on Jul. 17,2013, the disclosures of which are expressly incorporated by referenceherein.

TECHNICAL FIELD OF THE INVENTION

The present invention relates to an Ni-based alloy product, an Ni-basedalloy member produced of the Ni-based alloy product, a method forproducing the Ni-based alloy product, and a method for producing theNi-based alloy member.

DESCRIPTION OF BACKGROUND ART

How to improve thermal efficiency of high temperature devices, such asgas turbines and jet engines, is an important problem for many reasonsincluding the need to reduce environmental impacts. An effective way ofincreasing thermal efficiency is to increase service temperatures.

Currently, a turbine inlet temperature of about 1300° C. is standard ina gas turbine. On the other hand, turbine components applicable totemperatures around 1700° C. are becoming commercially practical. Also,for gas turbine components such as turbine blades, Ni-based alloys ofhigh heat-resistant superalloys are often used.

Meanwhile, high-strength Ni-based alloys applied to these gas turbines,jet engines, etc. derive their high mechanical strength fromprecipitating a γ′ phase (gamma prime phase, Ni₃Al) therein. A γ′ phaseis coherent with a γ phase in crystalline lattice, and the γ′ phasecoherently precipitated in the γ phase (hereinafter referred to as a“coherent γ′ phase”) contributes greatly to the improvement inmechanical strength. In other words, the mechanical strength of Ni-basedalloy members used in gas turbines, etc. can be improved by increasingthe amount of the precipitated γ′ phase. However, such high-strengthNi-based alloy members with a high content of the precipitated γ′ phasehave extremely poor cold workability due to their high hardness, andtherefore high-strength Ni-based alloy members are not usuallycold-worked.

For example, turbine blades mentioned above are produced of Ni-basedalloys by precision forging, in which a γ′ phase precipitate is presentat a ratio of 36 to 60 volume %, and cold working is not carried out inthe production process due to their high hardness.

On the other hand, as for combustor components produced by cold working,hardness can be reduced by using Ni-based alloys in which a γ′ phaseprecipitate is present at a controlled ratio of 30 volume % or lower,thereby making cold working possible. However, such combustor componentsand other articles that can be cold-worked have lower mechanicalstrength than turbine blades or the like produced of Ni-based alloysincluding a γ′ phase precipitate at a ratio of 36 to 60 volume %. And,such Ni-based alloys including a γ′ phase precipitate of 30 volume % orlower are not adequate to fully satisfy requirements for the capabilityto tolerate increasingly high temperatures, as mentioned above.

As seen from the above, what is strongly needed in the art is to developan Ni-based alloy member that is produced of an Ni-based alloy includinga γ′ phase precipitate of 36 to 60 volume % and having a high durabletemperature and that further has good cold workability. Also, a methodfor producing such a member is required.

Patent Literature 1 discloses a method for making an Ni-based superalloyarticle having a controlled grain size from a forging preform. In PatentLiterature 1, there is described a controlling method of a grain size ofan Ni-based superalloy, comprising the steps of hot die forging as theinitial forging operations and isothermal forging as the subsequentforging operations. With this controlling method, a uniform grain sizeof approximately ASTM 6 to 8 can be achieved by carrying out hot dieforging for the initial upset followed by isothermal forging and, ifnecessary, subsolvus annealing to provide a microstructure suitable forsupersolvus heat treatment. It also describes that the hot die forgingcauses partial or complete recrystallization of the microstructure,which facilitates superplastic deformation in the subsequent isothermalforging operations. Moreover, Examples disclosed in Patent Literature 1include a description about grain sizes when heat treatment is appliedat 1850° F., 1900° F., and 1925° F.

CITATION LIST Patent Literatures

Patent Literature 1: Japanese Unexamined Patent Application PublicationNo. Hei 9(1997)-302450.

SUMMARY OF THE INVENTION Problems to be Solved by the Invention

With the method for controlling the grain size of an Ni-based superalloydescribed in Patent Literature 1, a uniform grain size can be achieved,and in addition, superplastic deformation can be facilitated. However,this does not solve the above-mentioned problem, that is, does not makeit possible to provide an Ni-based alloy member including a γ′ phaseprecipitate at a ratio of 36 to 60 volume % and which has a high durabletemperature and also has good cold workability. Furthermore, PatentLiterature 1 does not provide a method for producing the Ni-based alloymember.

The present invention has been made in view of the above problems, andit is an objective to provide: an Ni-based alloy member in which a γ′phase precipitate is present at a ratio of 36 to 60 volume % and whichhas a high durable temperature and also has good cold workability; amethod for producing the member; an Ni-based alloy product to be used asa precursor of the Ni-based alloy member; and a method for producing theproduct.

Solution to Problems

According to one aspect of the present invention, there is provided anNi-based alloy product having a two-phase structure composed of a γphase and a γ′ phase that is incoherent with the γ phase in crystallinelattice parameters (hereinafter referred to as an “incoherent γ′phase”), in which the incoherent γ′ phase is present at a ratio of 20volume % or higher in the two-phase structure.

A hardness of the Ni-based alloy product can be decreased withincreasing contents of the incoherent γ′ phase, thereby facilitatingcold working. More preferable precipitation ratio of the incoherent γ′phase is 25 volume % or higher. Also, the hardness is preferably 400 Hvor lower, more preferably 370 Hv or lower.

Moreover, in order to enhance ductility in cold working and improve coldworkability, average grain size of the γ phase and the incoherent γ′phase is preferably 100 μm or smaller, more preferably 50 μm or smaller.

The same advantages of the invention can be obtained even when carbidesand different phases such as an η (eta) phase are present besides theincoherent γ′ phase. However, the total of such different phases ispreferably 15 volume % or less.

Furthermore, the advantages of the present invention can be obtainedeven when some precipitates of a fine-grained coherent γ′ phase arepresent in the γ phase. However, it is preferable that the amount of thecoherent γ′ phase be limited to a minimum.

The Ni-based alloy product according to the present invention isexcellent in cutting machinability as well as in cold workability.

In order to produce the Ni-based alloy product according to the presentinvention, hot forging needs to be performed in a temperature rangewhere the two phases of the γ phase and the incoherent γ′ phase cancoexist. The reason is not only to precipitate the incoherent γ′ phasebut also to obtain a fine microstructure by inhibiting the coarsening ofthe γ phase by the incoherent γ′ phase.

The hot forging needs to be performed at temperatures equal to or higherthan 1000° C., at which the mechanical strength of the incoherent γ′phase becomes lower. Furthermore, it is desirable that the incoherent γ′phase be present at a ratio of 10 volume % or higher during the hotforging.

After the forging, the hardness of the Ni-based alloy can be decreasedby increasing the incoherent γ′ phase, resulting in further enhanced hotworkability.

In order to increase the incoherent γ′ phase, it is effective to conducthomogenization heat treatment at a temperature equal to or higher than1000° C. and within a temperature range where the two phases of the γphase and the γ′ phase coexist, preferably at a heating temperature ofthe final forging. And, after the homogenization heat treatment, it iseffective to carry out slow cooling to a temperature 100° C. or morebelow the homogenization heat treatment temperature.

This slow cooling inhibits the precipitation of the coherent γ′ phaseinto the γ phase, which makes it possible to increase the incoherent γ′phase.

A cooling rate of 100° C./h or slower is effective; a cooling rate of50° C./h or slower is significantly effective; and a cooling rate of 20°C./h or slower is the most preferable.

Besides, an Ni-based alloy member according to the present invention isa Ni-based alloy member produced through cold working (including cuttingmachining), annealing, and solution and aging heat treatment of theNi-based alloy product described above. And, the Ni-based alloy membercomprises a γ phase and a coherent γ′ phase, in which the coherent γ′phase is present at a ratio of 36 to 60 volume %, and has apredetermined shape.

When conducting solution heat treatment to redissolve the incoherent γ′phase into a matrix, it is effective to apply a heat treatment attemperatures above a temperature at which the incoherent γ′ phasedissolves and becomes a solid solution completely. However, in the casewhere a grain size of the matrix becomes too coarse and the propertiesare degraded by the heat treatment, the coarsening of the crystallinegrains can be inhibited by applying the solution heat treatment attemperatures at which the incoherent γ′ phase remains to some extent. Inthis case, the amount of the residual incoherent γ′ phase is preferably10 volume % or less.

In addition, a method of an Ni-based alloy member according to thepresent invention includes the step of producing a precursor of anNi-based alloy member that has a predetermined shape by cold-working theNi-based alloy product produced by the method described above. Theprecursor of an Ni-based alloy member is subjected to solution and agingheat treatment so as to produce an Ni-based alloy member comprises a γphase and a coherent γ′ phase, wherein the coherent γ′ phase is presentat a ratio of 36 to 60 volume %.

Advantages of the Invention

According to an Ni-based alloy product and a method for producing theproduct of the present invention, the Ni-based alloy product produced byhot forging has a two-phase structure composed of a γ phase and a γ′phase that is incoherent with the γ phase, wherein the γ′ phase ispresent at a ratio of 20 volume % or higher, which leads to excellentcold workability in the Ni-based alloy product. Also, according to anNi-based alloy member and a method for producing the member of thepresent invention, by subjecting the above-mentioned Ni-based alloyproduct to cold working, forming it into a predetermined shape, and thensubjecting it to solution and aging heat treatment, there can beobtained an Ni-based alloy member having a high durable temperature, inwhich the Ni-based alloy member comprises a γ phase and a coherent γ′phase, the coherent γ′ phase being present at a ratio of 36 to 60 volume%.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a flowchart showing a method for producing an Ni-based alloymember according to a first embodiment of the present invention;

FIG. 2 is a schematic drawing showing a perspective view of an Ni-basedalloy product according to an embodiment of the present invention;

FIG. 3A is a schematic drawing showing a microstructure of an Ni-basedalloy product as a comparative example, FIG. 3B is a schematic drawingshowing a microstructure of an Ni-based alloy product after beingsubjected to hot forging as an inventive example, and FIG. 3C is aschematic drawing showing a microstructure of an Ni-based alloy memberobtained by subjecting a precursor of an Ni-based alloy member producedby cold-working the Ni-based alloy product of FIG. 3B to solution andaging heat treatment;

FIGS. 4A, 4B, and 4C each are a schematic drawings of an Ni-based alloymember according to an embodiment of the present invention;

FIG. 5 is a flowchart showing a method for producing an Ni-based alloymember according to a second embodiment of the present invention;

FIG. 6 is a graph showing test results that define an optimal range ofthe amount of a precipitated γ′ phase that is incoherent with a γ phasein a hot forged Ni-based alloy product; and

FIG. 7 is a graph showing a property ratio between a sample subjected tohot forging and solution and aging heat treatment and another samplesubjected to hot forging, cold working, and solution and aging heattreatment.

DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

Preferred embodiments of an Ni-based alloy product, a method forproducing the product, an Ni-based alloy member, and a method forproducing the member according to the present invention will bedescribed below with reference to the accompanying drawings.

(First Embodiment of Method for Producing Ni-Based Alloy Member)

FIG. 1 is a flowchart showing a method for producing an Ni-based alloymember according to a first embodiment of the present invention, andFIG. 2 is a schematic drawing showing a perspective view of an Ni-basedalloy product according to an embodiment of the present invention. Also,FIG. 3A is a schematic drawing showing a microstructure of an Ni-basedalloy product as a comparative example; FIG. 3B is a schematic drawingshowing a microstructure of an Ni-based alloy product after beingsubjected to hot forging as an inventive example, and FIG. 3C is aschematic drawing showing a microstructure of an Ni-based alloy memberobtained by subjecting a precursor of an Ni-based alloy member producedby cold-working the Ni-based alloy product of FIG. 3B to solution andaging heat treatment.

In the method for producing an Ni-based alloy member shown in theflowchart of FIG. 1, first, an Ni-based alloy product to be a basematerial for an Ni-based alloy member is produced, and then an Ni-basedalloy member is produced using this Ni-based alloy product.

An Ni-based alloy member produced by the production method according tothe present invention is made up of a γ phase and a γ′ phase that iscoherent with the γ phase, wherein the γ′ phase is present at a ratio of36 to 60 volume %, and has a high durable temperature. Morespecifically, the object to be produced by the production method of thepresent invention is an Ni-based alloy member wherein a γ′ phase that isthermodynamically stable in a temperature range of 700 to 900° C., inwhich the Ni-based alloy member is to be used, is present at a ratio of36 to 60 volume %.

In producing an Ni-based alloy member of such high mechanical strength,first, an Ni-based alloy product (a product as a production basematerial for the Ni-based alloy member) that has a two-phase structurecomposed of a γ phase and an incoherent γ′ phase, wherein the incoherentγ′ phase is present at a ratio of 20 volume % or higher, is produced byhot-forging an Ni-based alloy material at a temperature equal to orhigher than 1000° C. and at which the γ′ phase is precipitated at aratio of 10 volume % or higher (step S10 in FIG. 1). The Ni-based alloymaterial has an ingredient composition in which a γ′ phase at a ratio of36 to 60 volume % can be precipitated.

An example of the ingredient composition of the Ni-based alloy productwould be 12% of Co, 14% of Cr, 3.7% of Al, 2.6% of Ti, 1% of Nb, 1% ofW, 2% of Mo, 0.01% of C, and the balance of Ni (all in volume %),wherein an incoherent γ′ phase is present at a ratio of 20 volume % orhigher.

An Ni-based alloy product as an inventive example produced by hotforging has a microstructure shown in FIG. 3B.

In FIG. 3B, the γ phase M′ and the incoherent γ′ phase P′ are completelydifferent in crystal alignment, and their crystalline grains are locatedthrough the grain boundaries B of an incoherent interface. In otherwords, the incoherent γ′ phase P′ may be regarded as an excludedprecipitate from a crystalline grain of the γ phase M′.

Incidentally, in the γ phase M′, Ni and Al atoms are randomly arranged,but in the γ′ phase P′, Ni and Al atoms are regularly arranged. Whileboth are based on a face-centered cubic lattice, they are different asprecipitates.

For comparison with the microstructure of the Ni-based alloy product ofthe inventive example shown in FIGS. 3A and 3B are a schematic drawingshowing a microstructure of an Ni-based alloy product as a comparativeexample produced without being subjected to hot forging.

As shown in FIG. 3A, in the Ni-based alloy product produced withoutbeing subjected to hot forging, the γ′ phase P is precipitated as aninclusion in a circular shape (a substantially circular shape) withinthe crystalline grains of the γ phase M, and the crystalline grains ofthe γ phase M are adjacent to each other via the grain boundaries B.Since the γ phase M and the γ′ phase P are connected with each otherwithout the grain boundaries B, a coherent interface would be formed onthe interface between the two. In other words, this γ′ phase P can bereferred to as a coherent γ′ phase P.

Meanwhile, a γ′ phase generally has good lattice coherence with a γphase of a matrix. Therefore, a γ′ phase P precipitated within acrystalline grain of a γ phase M like FIG. 3A is coherent with the γphase M.

The inventors came up with a technical idea in which this γ′ phase P isnot significantly higher in mechanical strength than the γ phase M, andthat the coherent interface between the γ phase M and the γ′ phase Pwould enhance the mechanical strength of an Ni-based alloy member.

In other words, the inventors considered that the presence of a coherentinterface between a γ phase M and a γ′ phase P, as shown in FIG. 3A,results in poor cold workability of a high-strength Ni-based alloymember. Based on the above idea, the inventors have arrived at aninnovative technical idea in that the formation of a microstructurehaving no coherent interface between the γ phase and the γ′ phase at astage prior to cold working can lower the mechanical strength andhardness of the Ni-based alloy member temporarily at the stage of coldworking and thus improve its cold workability.

So, by carrying out hot forging or applying heat treatment after the hotforging at a temperature equal to or higher than 1000° C. and at whichtwo phases of a γ phase and a γ′ phase can coexist, there can beproduced an Ni-based alloy product having a two-phase structure in whicha γ phase M′ and a γ′ phase P′ that is incoherent with the γ phase M′are aligned via incoherent grain boundaries B as shown in FIG. 3B,instead of forming a coherent interface between a γ phase and a γ′ phaselike FIG. 3A. And then, by subjecting a relatively soft Ni-based alloyproduct to cold working, it is made possible to facilitate theproduction of an Ni-based alloy member of a desired shape.

Referring back to FIG. 1, a precursor of an Ni-based alloy member of adesired shape is produced by cold-working an Ni-based alloy product 1produced by hot forging (step S20).

Herein, “cold working” means working the Ni-based alloy product 1 intothe shape of a desired final Ni-based alloy member by, for example,forging, rolling, or molding at a room temperature.

Because the Ni-based alloy product 1 used has the microstructure shownin FIG. 3B and is relatively soft, it has low mechanical strength at aroom temperature and therefore exhibits excellent cold workability.

Enhancing ductility is effective in further improving this coldworkability, and it is preferable that the crystalline grains of boththe γ phase M′ and the incoherent γ′ phase P′ that form the Ni-basedalloy product 1 be adjusted to 100 μm or smaller in grain size. It ismore preferable that they be adjusted to 50 μm or smaller in grain size.

Regarding this grain size, the inventors have proved that by performingstep S10, namely, the step of hot-forging an Ni-based alloy basematerial at a temperature equal to or higher than 1000° C. and at whicha γ′ phase and a γ phase can coexist, a γ′ phase that is incoherent withthe γ phase is precipitated, and this precipitated γ′ phase inhibits thegrain growth of the γ phase. As a result, the grain size of both the γphase and the γ′ phase can be adjusted to 100 μm or smaller.

By this cold working, there is produced a precursor of an Ni-based alloymember that is a precursor of Ni-based alloy members such as plates,rod-shaped wires, and even turbine blades to be used as gas turbinecomponents.

However, the precursor of an Ni-based alloy member produced in step S20has a microstructure in which no coherent interface is present betweenthe γ phase and the γ′ phase to contribute to the enhancement ofmechanical strength. Therefore, the precursor itself is not suitable forapplication as high-strength members.

Then, the precursor of an Ni-based alloy member is subjected to solutionheat treatment so as to redissolve the incoherent γ′ phase into amatrix. Subsequently, the precursor is subjected to aging heat treatmentso as to precipitate a coherent γ′ phase as an inclusion in thecrystalline grains of the γ phase, which causes the formation of acoherent interface between the γ phase and the γ′ phase. Thus there isproduced an Ni-based alloy member that has the microstructure shown inFIG. 3C (see step S30).

Here, the microstructure shown in FIG. 3C contains a γ′ phase Pcoherently precipitated within a γ phase M as a matrix, and has acoherent interface formed between the γ phase M and the γ′ phase P,resulting in an Ni-based alloy member in which the γ′ phase P that isthermodynamically stable is present at a ratio of 36 to 60 volume %.

Examples of the Ni-based alloy member produced in step S30 are shown inFIGS. 4A to 4C. The Ni-based alloy member 10 shown in FIG. 4A is aplate, the Ni-based alloy member 10A shown in FIG. 4B is a wire, and theNi-based alloy member 10B shown in FIG. 4C is a turbine blade.

Each of these Ni-based alloy members 10, 10A and 10B contains a γ′ phaseat a ratio of 36 to 60 volume % or higher and has a high durabletemperature due to a coherent interface formed between a γ phase and aγ′ phase that is coherent with this γ phase.

As described above, according to the production flow shown in FIG. 1, anNi-based alloy member that has a high durable temperature and isexcellent in cold workability can be provided by the following steps:hot-forging a base material of a high-strength Ni-based alloy containinga γ′ phase precipitate in an amount of 36 volume % or larger to exercisestructure control to cause the precipitation of a γ′ phase that isincoherent with the γ phase so as to produce an Ni-based alloy productthat is relatively soft and excellent in cold workability; cold-workingthis Ni-based alloy product into a desired shape; and then subjecting itto solution and aging heat treatment to exercise structure control tocause the precipitation of a γ′ phase that is coherent with the γ phaseso as to produce a high-strength Ni-based alloy member. After the hotworking, the Ni-based alloy product may be reheated to the final forgingtemperature for homogenization and then air-cooled before the coldworking.

(Second Embodiment of Method for Producing Ni-Based Alloy Member)

FIG. 5 is a flowchart showing a method for producing an Ni-based alloymember according to a second embodiment of the present invention.

The production method for an Ni-based alloy member shown in FIG. 5 is aproduction method characterized in that it has an additional step ofsubjecting an Ni-based alloy product to heat treatment following thestep S10 in which the Ni-based alloy product is produced by hot forgingat a temperature equal to or higher than 1000° C. In this additionalstep, the Ni-based alloy product is subjected to homogenization heattreatment at a temperature equal to or higher than 1000° C. and at whichthe γ phase and the γ′ phase coexist, and slow-cooled to a temperature100° C. or more below the homogenization heat treatment temperatures(see step S10′). It is then cooled to a room temperature before beingsubjected to cold working.

For example, in the case where hot forging is performed at temperaturesaround 1200° C. in the initial stage and at around 1150° C. in the finalstage, the subsequent heat treatment is applied for a predetermined timeat a temperature around 1100° C., which is below the final stagetemperature of the hot forging of about 1150° C., and then heattreatment is applied while controlling the temperature by slow-coolingthe Ni-based alloy product to temperatures around 1000° C. or 900° C.

The inventors have revealed that by applying heat treatment after hotforging for a predetermined time at a temperature below the hot forgingtemperatures in the way described above, the incoherent γ′ phase can beincreased to further lower the hardness of the Ni-based alloy product,which results in further improved cold workability.

[Cold Workability Verification Tests and Results Thereof]

The inventors produced test pieces of different ingredient compositionsunder different production conditions and conducted tests to verify thecold workability of each test piece. Table 1 below shows the ingredientcompositions of the test pieces, and Table 2 shows the productionconditions of the test pieces and cold working test results. Also, asfor the test pieces for which heat treatment was applied after hotforging during their production, the details of the heat treatments A, Band C in Table 2 are shown in Table 3.

TABLE 1 Ingredient Compositions of Test Pieces (vol. %). Test No. Ni CrCo Mo W Ti Al C B Zr Nb Fe Others Comparative Balance 16 15 3 1.3 4 2.80.025 0.018 0.03 0 0 Example 1 Comparative Balance 16 15 3 1 5 2.5 0.0250.018 0.03 0 0 Example 2 Comparative Balance 13.5 20 2.8 1.2 5.8 2.30.015 0.015 0.03 0 0 Example 3 Comparative Balance 13.5 20 2.8 1.2 4.8 30.015 0.015 0.03 0 0 Example 4 Comparative Balance 16 5 4 3 4 2.7 0.010.001 0.003 0 0 Example 5 Comparative Balance 16 15 3 1.3 4.9 2.5 0.0250.001 0.003 0 0 Example 6 Inventive Balance 13 0 5 0 5 2.7 0.002 0.0180.04 0 0 Example 1 Inventive Balance 16 10 0 4 3 3.6 0.001 0.009 0 0 5Example 2 Inventive Balance 17 10 2 1 3 3.8 0.02 0.001 0.001 2 0 1.0TaExample 3 Inventive Balance 16 7 4 1 4 2.7 0.006 0.001 0.003 0 0 1.0TaExample 4 Inventive Balance 16 7 4 1 0.5 5 0.006 0.001 0.003 0.8 0 0.5HfExample 5 Inventive Balance 14 12 2 1 2.6 3.7 0.01 0.012 0.04 1 0Example 6 Inventive Balance 18 26 0 0 1.8 4 0.04 0.02 0.02 2.2 2 Example7 Inventive Balance 16 5 4 3 4 2.7 0.01 0.001 0.003 0 0 Example 8Inventive Balance 16 15 3 1.3 4.9 2.5 0.025 0.001 0.003 0 0 Example 9Inventive Balance 15.7 8.5 3.1 2.7 3.4 2.3 0.015 0.01 0.03 1.1 4 Example10

TABLE 2 Production Conditions of Test Pieces and Cold Working TestResults. Amount of γ′ Phase at Amount Hardness Service Hot Forging HotForging of before Temperature Start End Heat Treatment Incoherent ColdCold (700° C.) Temperature Temperature after γ′ Phase Working WorkingTest No. (vol. %) (° C.) (° C.) Hot Forging (vol. %) (Hv) Test ResultComparative 42 Not performed — — 480 NG Example 1 Comparative 45 11801180 Not performed 0 470 NG Example 2 Comparative 46 1180 1180 Notperformed 0 466 NG Example 3 Comparative 47 1165 1180 Not performed 7455 NG Example 4 Comparative 42 1180 1180 Not performed 0 470 NG Example5 Comparative 44 1180 1150 Not performed 0 460 NG Example 6 Inventive 461180 1050 Not performed 20 365 OK Example 1 Inventive 47 1180 1000 Notperformed 30 360 OK Example 2 Inventive 45 1180 1050 Not performed 21390 OK Example 3 Inventive 43 1180 1000 Not performed 27 330 OK Example4 Inventive 47 1180 1150 Heat treatment A 30 365 OK Example 5 Inventive46 1150 1150 Heat treatment B 29 360 OK Example 6 Inventive 43 1180 1150Heat treatment C 35 320 OK Example 7 Inventive 42 1180 1150 Heattreatment A 30 340 OK Example 8 Inventive 44 1180 1150 Heat treatment B32 325 OK Example 9 Inventive 37 1180 1120 Heat treatment B 22 355 OKExample 10

TABLE 3 Heat Treatment A Held at 1100° C. for 1 hour, then cooled to1000° C. at rate of 10° C./hour, and then water-cooled Heat Treatment BHeld at 1100° C. for 1 hour, then cooled to 1000° C. at rate of 50°C./hour, then cooled to 950° C. at rate of 20° C./hour, and thenair-cooled Heat Treatment C Held at 1100° C. for 1 hour, then cooled to900° C. at rate of 5° C./hour, and then air-cooled

In producing each test piece, the base material of 20 kg was melted byvacuum induction melting, subjected to homogenization heat treatment,and subsequently hot-forged under the conditions shown in Table 2 into around bar with a diameter of 15 mm.

In Comparative Example 1, hot forging was not performed, whereas inComparative Examples 2 to 6, hot forging was performed. Hot forging wasperformed also in Inventive Examples 1 to 10, and as for InventiveExamples 5 to 10, one of the heat treatments A to C shown in Table 3 wasapplied after the hot forging.

The microstructure of each test piece was observed after the hot forgingor after the subsequent heat treatment, and the content ratios of the γphase and the incoherent γ′ phase were measured.

Furthermore, the cold working tests were conducted in the followingprocedure. First, each obtained round bar with a diameter of 15 mm wasreduced in diameter 1 mm by 1 mm, by cold drawing. The cold drawing wasperformed three times until the diameter was reduced to 12 mm.

The cold working test results for the test pieces that could not bedrawn successfully are denoted as “NG” in Table 2.

In contrast, the cold working test results for the test pieces thatcould be drawn successfully into a test piece with a diameter of 13 mmwithout cracking are denoted as “OK” in Table 2. Some test pieces weresubsequently subjected to annealing at temperatures between 1000° C. and1100° C. and cold working repeatedly to be successfully worked into awire rod with a diameter of 3 mm.

As shown in Table 2, the cold working test results for test pieces ofComparative Examples 1 to 6 were all “NG”, whereas the cold working testresults for test pieces of Inventive Examples 1 to 10 were all “OK”. Inparticular, it was easy to perform cold working of test piecescontaining an incoherent γ′ phase precipitate in an amount of 25% orlarger and having a hardness of 370 Hv or lower.

As for test pieces of Comparative Examples 1 to 6, despite the hotforging performed, the amount of the incoherent γ′ phase remained 0volume %, resulting in a Vickers hardness (Hv) of over 400 before coldworking, with which cold working was impossible. This is because, exceptfor Comparative Example 4, the hot forging temperatures were higher thanthe solvus temperature of the γ′ phase and therefore no γ′ phaseprecipitation occurred during the hot forging. In Comparative Example 4,the hot forging temperatures were slightly lower than the solvustemperature of the γ′ phase, and therefore an incoherent γ′ phase wasprecipitated in a small amount, which, however, was not enough toimprove cold workability. The solvus temperatures of the γ′ phase ofComparative Examples 1 to 6 were 1134° C., 1157° C., 1183° C., 1173° C.,1115° C., and 1154° C., respectively.

In contrast, the Vickers hardness (Hv) of each test piece of InventiveExamples 1 to 10 was lower than 400, which permits cold working.

In particular, Inventive Examples 5 to 10, for which any one of the heattreatments A to C was applied after the hot forging, each exhibited aVickers hardness (Hv) that was relatively low as compared with InventiveExamples 1 to 3, for which no heat treatment was applied after the hotforging.

As can be seen from the above, it has been demonstrated that thehardness of an Ni-based alloy product can be further lowered to furtherimprove its cold workability by applying homogenization heat treatmentat a temperature equal to or higher than 1000° C. and within atemperature range in which the γ phase and the γ′ phase coexist afterperforming hot forging in the way described above and subsequentlyperforming slow cooling to a temperature 100° C. or more below thehomogenization heat treatment temperature.

Incidentally, test pieces of Inventive Examples 1 to 8 were successfullyworked into a wire with a diameter of 2 mm by being subjected toannealing and cold drawing repeatedly after the first cold working test.

A relationship between the amount of the precipitated incoherent γ′phase and the Vickers hardness before the cold working in Table 2 isshown in a graph form in FIG. 6.

FIG. 6 teaches that the amount of precipitation of the incoherent γ′phase to the γ phase meets an inflection point at 20 volume %, and thatthe Vickers hardness greatly decreases in a range of the amount equal toor larger than 20 volume %. It also teaches that in this range of theamount equal to or larger than 20 volume %, the Vickers hardness islower than 400 Hv, which indicates that cold working is possible. Basedon these results, it has been determined that the amount of theprecipitated incoherent γ′ phase contained in an Ni-based alloy productproduced by hot forging at a temperature equal to or higher than 1000°C. is defined to be 20 volume % or larger.

FIG. 7 is a graph showing a property ratio between a sample subjected tohot forging and solution and aging heat treatment and another samplesubjected to hot forging, cold working, and solution and aging heattreatment.

Here, tensile testing was conducted in two cases, at a room temperatureand at 700° C. Also, creep testing was conducted at 700° C. and a loadstress of 350 MPa.

FIG. 7 teaches that the two test pieces exhibit almost the same tensileproperty and creep property. Therefore, it has been found that anNi-based alloy member produced by being subjected to hot forgingfollowed by cold working and subsequently to solution and aging heattreatment as with the production method according to the presentinvention has a mechanical strength equivalent to that of anotherNi-based alloy member produced by a production method in which coldworking is not performed.

While the preferred embodiments of the present invention have beendescribed above with reference to the accompanying drawings, it shouldbe noted that the specific constitution is not to be construed aslimited to the embodiments and that any design modifications, etc. madewithout departing from the spirit and scope of the present invention areto be included in the present invention.

LEGEND

1 . . . Ni-based alloy product; 10, 10A, 10B . . . Ni-based alloymember; B . . . grain boundary; M . . . γ phase (matrix); P . . . γ′phase (γ′ phase coherent with γ phase); and P′ . . . γ′ phase (γ′ phaseincoherent with γ phase).

1. A method for producing an Ni-based alloy product having a two-phasestructure composed of crystalline grains of a γ phase and crystallinegrains of the γ′ phase, in which the γ′ phase is an incoherent γ′ phasebeing located through γ phase grain boundaries of an incoherentinterface, and in which the incoherent γ′ phase is present at a ratio of20 volume % or higher in the two-phase structure, the method comprising:the step of melting and casting a first Ni-based alloy material toproduce a second Ni-based alloy material; and the step of hot-forgingthe second Ni-based alloy material obtained through the melting andcasting step at temperatures equal to or higher than 1000° C. andwherein the γ and the γ′ phases coexist to produce a third Ni-basedalloy material having an ingredient composition in which a γ′ phase at aratio of 36 to 60 volume % can be precipitated in a temperature range of700 to 900° C.; and the step of forming the Ni-based alloy product fromthe third Ni-based alloy material.
 2. The method for producing anNi-based alloy product according to claim 1, wherein the γ′ phaseprecipitates at a ratio of 10 volume % or higher in the hot-forgingstep.
 3. The method for producing an Ni-based alloy product according toclaim 1, wherein the method further comprises the step of subjecting thethird Ni-based alloy material obtained through the hot-forging step to ahomogenization heat treatment prior to forming the Ni-based alloyproduct by (1) holding a temperature equal to or higher than 1000° C.but equal to or lower than the hot-forging temperature and wherein the γand the γ′ phases coexist, and (2) slow-cooling to a temperature by 100°C. or more below the holding temperature.
 4. The method for producing anNi-based alloy product according to claim 2, wherein the method furthercomprises the step of subjecting the third Ni-based alloy materialobtained through the hot-forging step to a homogenization heat treatmentprior to forming the Ni-based alloy product by (1) holding a temperatureequal to or higher than 1000° C. but equal to or lower than thehot-forging temperature and wherein the γ and the γ′ phases coexist, and(2) slow-cooling to a temperature by 100° C. or more below the holdingtemperature.
 5. The method for producing an Ni-based alloy productaccording to claim 3, wherein a rate of the slow-cooling in thehomogenization heat treatment is equal to or slower than 100° C./h. 6.The method for producing an Ni-based alloy product according to claim 5,wherein a rate of the slow-cooling in the homogenization heat treatmentis 5° C./h or more and 50° C./h or less.
 7. The method for producing anNi-based alloy product according to claim 4, wherein a rate of theslow-cooling in the homogenization heat treatment is equal to or slowerthan 100° C./h.
 8. The method for producing an Ni-based alloy productaccording to claim 7, wherein a rate of the slow-cooling in thehomogenization heat treatment is 5° C./h or more and 50° C./h or less.9. The method for producing an Ni-based alloy product according to claim1, wherein during the hot-forging step, the temperature of thehot-forging is lowered to a temperature by 100° C. or more below ahot-forging start temperature.
 10. The method for producing an Ni-basedalloy product according to claim 2, wherein during the hot-forging step,the temperature of the hot-forging is lowered to a temperature by 100°C. or more below a hot-forging start temperature.